High-strength steel sheet with excellent crashworthiness characteristics and formability and method of manufacturing the same

ABSTRACT

Provided is a high-strength steel sheet including, in % weight, carbon (C): 0.04 to 0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.8 to 2.5%, molybdenum (Mo): 0.15% or less (excluding 0%), chromium (Cr): 1.0% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.01% or less, aluminum (Al): 0.01 to 0.5%, nitrogen (N): 0.01% or less, boron (B): 0.01% or less (excluding 0%), antimony (Sb): 0.05% or less (excluding 0%), one or more of titanium (Ti): 0.003 to 0.06% and niobium (Nb): 0.003 to 0.06%, a balance of Fe and other unavoidable impurities, and contents of the C, the Si, the Al, the Mo and the Cr satisfy the following Expression 1: Expression 1: {(2×(Si+Al))+Mo+Cr}/C≥15. The high-strength steel sheet comprises: a ferrite phase, a bainite phase, a martensite phase, and a residual austenite phase, the ferrite phase being less than 40% of area fraction in the microstructure.

TECHNICAL FIELD

The present invention relates to a high-strength steel sheet used for anautomotive structural member, and more particularly, to a high-strengthsteel sheet with excellent crashworthiness characteristics andformability and a method for manufacturing the same.

BACKGROUND ART

In the automotive industry, regulations on carbon dioxide emissions aswell as environmental and safety regulations have become increasinglysevere, and fuel regulations for automobiles are also being tightened.In addition, the US Highway Safety Insurance Association is increasinglyregulating collision safety to protect passengers, and requires severecrash performance, such as a 25% overlap from 2013 regulations.

Research to reduce the weight of automobiles as a means of solving theproblems in the automobile industry is being carried out continuously.

Generally, high-strength automobile materials can be classified into aprecipitation hardening steel, a bake hardenable steel, a solid solutionstrengthening steel, a transformed hardened steel, etc.

In the transformed hardened steel among them, dual phase steel (DPsteel), transformation induced plasticity steel (TRIP steel), complexphase steel (CP steel) and the like are included. Such transformedhardened steel is called advanced high strength steel (AHSS).

DP steel is a steel in which hard martensite is finely and homogeneouslydispersed in soft ferrite to ensure high strength. CP steel contains twoor three phases of ferrite, martensite and bainite, and includesprecipitation hardening elements such as Ti, Nb and the like to enhancestrength. TRIP steel is a type of steel capable of securing highstrength and high ductility, causing martensite transformation whenremaining fine austenite, which is homogeneously dispersed, is processedat room temperature.

Meanwhile, recently, a steel sheet for automobiles has required evenhigher strength to improve fuel economy or durability. An ultrahigh-strength steel sheet having a tensile strength of 980 MPa or moreis increasingly used for automobile body structures or a reinforcingmaterial in aspects of collision safety and passenger protection.

Particularly, a high strength steel having excellent yield strength isused in a structural member such as a member, a seat rail, a pillar andthe like for enhancing the crash resistance of the vehicle body. Such astructural member has characteristics favorable to impact energyabsorption as the yield strength (YS) versus the tensile strength (TS),that is, the yield ratio (YS/TS), is high.

However, since ductility is usually reduced as the strength of the steelsheet is increased, there arises a problem is which formingprocessability is lowered.

Therefore, in order to secure both collision stability and componentformability, it is necessary to develop a material having high yieldstrength and excellent ductility. Further, since most of the parts to beprocessed are subjected to bending, the application of a steel materialhaving excellent bending performance is also required.

Therefore, in order to improve the crashworthiness characteristics andformability of high-strength steels, it is necessary to have low yieldstrength, improve bending performance and have excellent ductility likeDP steel, the most widely used among the transformation-strengtheninghigh strength steels, and the application of the high-strength steelsmay expand in various fields.

Meanwhile, water cooling is used during continuous annealing as atypical method for increasing yield strength. That is, steel may becracked in the annealing process, immersed in water and tempered toobtain a steel sheet having a tempered martensite structure in whichmartensite is tempered as a microstructure.

As a conventional technique related thereto, Patent Document 1 disclosesa method of continuously annealing a steel material containing carbon at0.18% or more, and water-cooling to room temperature, and performingoveraging treatment at a temperature of 120 to 300° C. for 1 to 15minutes to obtain a volume fraction of martensite of 80 to 97%.

However, although this technique can increase a yield ratio, the shapequality of the coil is poor due to the temperature deviation in thewidth direction and the longitudinal direction of the steel sheet duringwater cooling, causing defects such as deterioration in workability suchas the occurrence of cracks and the variation in material propertiesbetween different positions during forming.

As a technique of improving formability in high-tensile-steel sheets,Patent Document 2 discloses a steel sheet having a complex phase mainlycontaining martensite. That is, a method of producing a high-tensilesteel sheet in which fine precipitated copper particles having aparticle diameter of 1 to 100 nm are dispersed in a structure forimproving processability is proposed.

However, this technique requires excessive addition of Cu in an amountof 2 to 5 wt % in order to precipitate fine Cu particles, and thus hotshortness due to Cu may occur, and manufacturing costs may riseexcessively.

As still another example, Patent Document 3 discloses a steel sheethaving a microstructure containing ferrite as a matrix and containing 2to 10% by area of pearlite, and in which precipitation is strengthenedby adding elements such as Nb, Ti, V or the like which areprecipitation-strengthening elements, and having strength improved bygrain refinement.

In this case, although the hole expandability of the steel sheet isgood, there is a limit in increasing tensile strength, yield strength ishigh and ductility is low, which causes defects such as cracks duringpress forming.

As another example, Patent Document 4 discloses a cold-rolled steelsheet having high strength and high ductility simultaneously byutilizing a tempered martensite phase and also has a good plate shapeafter continuous annealing.

However, in this case, there is a problem in which the content of carbon(C) is as high as 0.2% or more, resulting in poor weldability andfurnace dent defect due to the addition of a large amount of Si.

(Patent Document 1) Japanese Patent Laid-Open Publication No.1992-289120(Patent Document 2) Japanese Patent Laid-Open Publication No.2005-264176(Patent Document 3) Korean Patent Laid-Open Publication No.2015-0073844(Patent Document 4) Japanese Patent Laid-Open Publication No.2010-090432SUMMARY OF THE INVENTION

An aspect of the present invention provides a high-strength steel sheethaving a tensile strength of 980 MPa or more, and more specifically,provides a high-strength steel sheet having low yield strength, highbending performance, excellent ductility and improved formability.

An aspect of the present invention provides a high-strength steel sheetwith excellent crashworthiness characteristics and formability,including, in % by weight, carbon (C): 0.04 to 0.15%, silicon (Si): 0.01to 1.0%, manganese (Mn): 1.8 to 2.5%, molybdenum (Mo): 0.15% or less(excluding 0%), chromium (Cr): 1.0% or less, phosphorus (P): 0.1% orless, sulfur (S): 0.01% or less, aluminum (Al): 0.01 to 0.5%, nitrogen(N): 0.01% or less, boron (B): 0.01% or less (excluding 0%), antimony(Sb): 0.05% or less (excluding 0%), one or more of titanium (Ti): 0.003to 0.06% and niobium (Nb): 0.003 to 0.06%, a balance as Fe and otherunavoidable impurities, and the relationship of C, Si, Al, Mo and Crsatisfies the following Expression 1, the high-strength steel sheetincluding a less than 40% area fraction of ferrite and the remainder asbainite, martensite and residual austenite, as a microstructure, whereinan area ratio (Fn/Ft) of non-recrystallized ferrite (Fn) in the totalferrite fraction (Ft) is 20% or less.

{(2×(Si+Al))+Mo+Cr}/C≥15   Expression 1

(Here, each element refers to a weight content.)

Another aspect of the present invention provides a method of producing ahigh-strength steel sheet with excellent crashworthiness characteristicsand formability, comprising: reheating a steel slab satisfying the alloycomposition and component relation (Expression 1) in a temperature rangeof 1050 to 1300° C., finishing hot-rolling the heated steel slab at atemperature equal to or higher than Ar3 to produce a hot-rolled steelsheet; coiling the hot-rolled steel sheet in a temperature range of 400to 700° C.; producing a cold-rolled steel sheet by cold-rolling afterthe coiling; continuously annealing the cold-rolled steel sheet in atemperature range of Ac1+30° C. to Ac3−20° C.; primary cooling at acooling rate of 10° C./s or less (excluding 0° C./s) up to 630 to 670°C. after the continuous annealing; secondary cooling at a cooling rateof 5° C./s or more up to 400 to 550° C. in hydrogen cooling equipmentafter the primary cooling; maintaining the temperature for 50 to 500seconds after the secondary cooling; hot-dip galvanizing after themaintaining; and final-cooling at a cooling rate of 1° C./s or higher toMs or lower after the hot-dip galvanizing, wherein the cold rolling iscarried out at a total reduction ratio of 30% or more, and each of thereduction ratio of the first to fourth stands is 15% or more.

Advantageous Effects

According to the present invention, a high-strength steel sheet havingimproved crashworthiness characteristics and formability can be providedby optimizing the alloy composition and the manufacturing conditions.

Particularly, since impact resistance is excellent due to high yieldstrength, and excellent ductility and bending performance can preventmachining defects such as cracks during press forming, it can besuitably applied to parts or the like requiring processing into acomplicated shape.

DESCRIPTION OF DRAWINGS

FIG. 1 shows a change in the interphase hardness ratio[(H_(B)+H_(M))/(2×H_(F))] according to a concentration ratio(corresponding to Expression 2) of Si, Al, Mo, Cr and C in a ferritephase in an embodiment of the present invention.

FIG. 2 shows a change in the product of yield strength and elongation(YS×El) according to the interphase hardness ratio[(H_(B)+H_(M))/(2×H_(F))] in an embodiment of the present invention.

FIG. 3 shows a change in a three-point bending angle according to theinterphase hardness ratio [(H_(B)+H_(M))/(2×H_(F))] in an embodiment ofthe present invention.

FIG. 4 shows a change in a three-point bending angle according to theproduct of yield strength and elongation (YS×El) in an embodiment of thepresent invention.

DETAILED DESCRIPTION OF THE INVENTION

The inventors of the present invention have conducted intensive studiesto develop materials having excellent crashworthiness characteristicswithout causing defects such as cracks during processing into automotivematerials having complex shapes.

As a result, it was confirmed that a high-strength steel sheet having astructure favorable for securing target physical properties can beprovided by optimizing the alloy composition and the manufacturingconditions, and thereby the present invention has been accomplished.

In particular, the present invention has technical significance in thatferrite is included as a microstructure, and the concentration ratio ofthe solid solution strengthening element in the ferrite is increased toinduce an increase in yield strength, the hardness ratio of bainite ormartensite phase in a hard phase is reduced to improve bendingprocessability.

In addition, the effect of inhibiting the grain growth can be obtainedby the solid solution strengthening elements in the ferrite, and thuseach phase is finely distributed, thereby relieving the local stressconcentration and greatly improving ductility.

Hereinafter, the present invention will be described in detail.

Preferably, a high-strength steel sheet with excellent crashworthinesscharacteristics and formability according to an aspect of the presentinvention includes, in % by weight, carbon (C): 0.04 to 0.15%, silicon(Si): 0.01 to 1.0%, manganese (Mn): 1.8 to 2.5%, molybdenum (Mo): 0.15%or less (excluding 0%), chromium (Cr): 1.0% or less, phosphorus (P):0.1% or less, sulfur (S): 0.01% or less, aluminum (Al): 0.01 to 0.5%,nitrogen (N): 0.01% or less, boron (B): 0.01% or less (excluding 0%),antimony (Sb): 0.05% or less (excluding 0%), one or more of titanium(Ti): 0.003 to 0.06% and niobium (Nb): 0.003 to 0.06%.

Hereinafter, the reason why the alloy composition of the high-strengthsteel sheet is controlled as described above will be described indetail. Here, unless otherwise specified, the content of each alloycomposition is % by weight.

C: 0.04 to 0.15%

Carbon (C) is the main element added for strengthening the transformedstructure of steel. The C improves the strength of the steel andpromotes the formation of martensite in the complex phase steel. As theC content increases, the amount of martensite in the steel increases.

However, when the content of C exceeds 0.15%, the strength is increaseddue to an increase in the amount of martensite in the steel, but adifference in strength from ferrite having a relatively low carbonconcentration is increased. Such a difference in strength causes aproblem in which ductility and the work hardening rate are loweredbecause the fracture occurs easily at the interface between phasesduring the addition of stress. In addition, there is a problem in whichwelding defects are generated in the processing of the parts of thecustomer due to low weldability. On the other hand, when the content ofC is less than 0.04%, it becomes difficult to secure desired strength.

Therefore, in the present invention, the content of C is preferablycontrolled to 0.04 to 0.15%, and more preferably 0.06 to 0.13%.

Si: 0.01 to 1.0%

Silicon (Si) is an element that stabilizes ferrite, and contributes tothe formation of martensite by promoting ferrite transformation andpromoting C concentration in untransformed austenite. In addition, it isan element which is effective for enhancing the strength of ferrite andreducing a phase hardness difference due to the high solid solutionstrengthening effect, and is effective for securing strength withoutlowering the ductility of the steel sheet.

For the above-mentioned effect, Si may be contained in an amount of0.01% or more. However, when the content exceeds 1.0%, surface scaledefects are caused, and the quality of the plated surface is poor andchemical treatment performance is deteriorated.

Therefore, in the present invention, it is preferable to control the Sicontent to 0.01 to 1.0%, and more preferably, the Si content may be inthe range of 0.1 to 0.8%.

Mn: 1.8 to 2.5%

Manganese (Mn) has the effect of refining the particles withoutdeterioration of ductility and precipitating sulfur (S) in steel as MnSto prevent hot shortness due to the formation of FeS. Further, the Mn isan element that strengthens the steel and serves to lower the criticalcooling rate at which a martensite phase is obtained in complex phasesteel, and is useful for forming martensite more easily.

When the content of Mn is less than 1.8%, not only can theabove-mentioned effect not be obtained, but also it is also difficult tosecure strength at the target level. On the other hand, when the contentexceeds 2.5%, there is a high possibility that problems in weldability,hot rollability and the like are likely to occur, martensite is formedexcessively so that the material is unstable, and a Mn-Band (Mn oxideband) is formed in the structure to increase the risk of the occurrenceof processing cracks and plate breakage. Further, there is a problem inwhich Mn oxide is eluted on the surface during annealing and greatlydeteriorates plating properties.

Therefore, in the present invention, it is preferable to control the Mncontent to 1.8 to 2.5%, and more preferably, the Mn content may be inthe range of 2.0 to 2.4%.

Mo: 0.15% or less (excluding 0%)

Molybdenum (Mo) is an element added to improve refinement of ferrite andstrength while retarding transformation of austenite into pearlite. SuchMo may improve the hardenability of steel to finely form martensite ingrain boundaries to enable the control of a yield ratio. However, sinceMo is an expensive element, there is a disadvantage in terms ofproduction of steel as the content of Mo is increased, and thus it ispreferable to suitably control the content of Mo.

In order to fully obtain the above-described effect, Mo may be added ina maximum amount of 0.15%. When the content of Mo exceeds 0.15%, thecost of an alloy is rapidly increased and thus economic efficiency islowered. Further, the ductility of steel may be deteriorated due toexcessive grain refinement and solid solution strengthening effects.

Therefore, in the present invention, it is preferable to control thecontent of Mo to be 0.15% or less, but the content is not 0%.

Cr: 1.0% or less (excluding 0%)

Chromium (Cr) is an element added to improve the hardenability of steeland ensure high strength. Such Cr is effective for forming martensiteand minimizes a decrease in ductility compared to an increase instrength, and thus it is advantageous for producing a complex phasesteel having high ductility. Especially, it is a solid solutionstrengthening element contributing to the enhancement of the strength offerrite.

In an aspect of the present invention, when the content of Cr exceeds1.0%, the effect is saturated and hot rolling strength is excessivelyincreased, thereby causing the cold rolling property to be poor.Further, there is a problem in which the fraction of the Cr-basedcarbide is increased and coarsened, and the size of the martensite iscoarsened after annealing, thereby leading to a decrease in elongation.

Therefore, in the present invention, it is preferable to control the Crcontent to be 1.0% or less, but the content is not 0%.

P: 0.1% or less

Phosphorus (P) is a substitutional element having the highest solidsolution-strengthening effect, and is an element favorable for securingstrength without improving in-plane anisotropy and greatly reducingformability. However, when such P is added excessively, the possibilityof the occurrence of brittle fracturing greatly increases, so that thepossibility of occurrence of plate breakage of the slab during hotrolling is increased, and the plated surface properties aredeteriorated.

Therefore, in the present invention, it is preferable to control the Pcontent to be 0.1% or less, but is not 0% because of the amountinevitably added.

S: 0.01% or less

Sulfur (S) is an element which is inevitably added as an impurityelement in steel, and deteriorates ductility and weldability, and thusit is preferable to control the content to be as low as possible.Particularly, since S has a problem of increasing the possibility ofgenerating hot shortness, it is preferable to control the content to0.01% or less, but the content is not 0% because of the amountinevitably added.

Al: 0.01 to 0.5%

Aluminum (Al) is an element added for finer grain size and deoxidationof steel. Further, Al is a ferrite-stabilizing element, and is effectivein distributing carbon in ferrite to austenite to improve thehardenability of martensite and is also effective in improving theductility of the steel sheet by effectively suppressing theprecipitation of carbides in bainite when being maintained in thebainite region.

It may be contained at 0.01% or more for the above-mentioned effect.When the content of Al exceeds 0.5%, it is advantageous in terms ofstrength improvement due to the grain refinement effect, but excessamounts of inclusions are formed during continuous casting insteelmaking, thereby increasing the possibility of surface defects inthe plated steel sheet. Further, an increase in manufacturing costs iscaused.

Therefore, in the present invention, the content of Al is preferablycontrolled to 0.01 to 0.5% or less.

N: 0.01% or less

Nitrogen (N) is an effective element for stabilizing austenite. However,when the content exceeds 0.01%, steel refining costs increase sharply,and the risk of cracking during continuous casting is greatly increaseddue to the formation of AIN precipitates.

Therefore, in the present invention, it is preferable to control the Ncontent to be 0.01% or less, but the content is not 0% because of theamount inevitably added.

B: 0.01% or less (excluding 0%)

Boron (B) is an advantageous element for retarding transformation ofaustenite into pearlite in the process of cooling during annealing.Further, it is also a hardenable element that inhibits the formation offerrite and promotes the formation of martensite.

When the content of B exceeds 0.01%, there is a problem in which anexcess amount of B is concentrated on the surface, resulting indeterioration of plating adhesion.

Therefore, in the present invention, it is preferable to control thecontent of B to 0.01% or less, but the content is not 0%.

Sb: 0.05% or less (excluding 0%)

Antimony (Sb) is distributed in grain boundaries and serves to retardthe diffusion of oxidizing elements such as Mn, Si, Al and the likethrough grain boundaries. Accordingly, the surface concentration of theoxide is suppressed and there is an advantageous effect in suppressingthe coarsening of the surface agglomerates due to the temperature riseand a change in the hot rolling process.

When the content of Sb is more than 0.05%, the effect is saturated,manufacturing costs are increased, and processability is lowered.

Therefore, in the present invention, it is preferable to control thecontent of Sb to 0.05% or less, but the content is not 0%.

One or more of Titanium (Ti): 0.003 to 0.06% and niobium (Nb): 0.003 to0.06%

Titanium (Ti) and niobium (Nb) are effective elements for increasingstrength and grain refinement by forming fine precipitates.Specifically, the Ti and Nb bond with C in steel to form nanoscale fineprecipitates, which serve to strengthen the matrix and decrease thephase hardness difference.

When the contents of Ti and Nb are less than 0.003%, respectively, theabove-mentioned effect cannot be sufficiently ensured. On the otherhand, when the contents of Ti and Nb are more than 0.06%, respectively,manufacturing costs increase and the precipitate is formed excessively,resulting in a decrease in ductility.

Therefore, in the present invention, the contents of Ti and Nb each arepreferably controlled to 0.003 to 0.06%.

The remainder of the present invention is iron (Fe). However,impurities, which are not intended, may be inevitably added from the rawmaterial or the surrounding environment in the common manufacturingprocess, and thus cannot be excluded. These impurities are notspecifically mentioned in this specification, as they are commonly knownto those skilled in the art.

Meanwhile, in order to have low yield strength as well as target highstrength and to improve bending performance and ductility to ensureexcellent formability in the present invention, the microstructure ofthe steel sheet is necessarily configured as follows.

Specifically, the high-strength steel sheet of the present invention mayinclude less than 40% (excluding 0%) by area fraction of ferrite and theremainder as bainite, martensite and residual austenite, as amicrostructure.

It is important to control the phase and the fraction for the complexphase steel including the soft phase and the hard phase combined tosatisfy a low yield ratio and high ductility while securing excellentbending properties. It should be noted that, unless otherwise specified,the fraction of the phase is area % in the present invention.

In the present invention, the ferrite phase is contained at less than40%, preferably, at 25% or more and less than 40%. Here, it ispreferable that the area ratio (Fn/Ft) of the non-recrystallized ferrite(Fn) in the total fraction of ferrite (Ft) is 20% or less (including0%). Here, when the area ratio (Fn/Ft) of the non-recrystallized ferriteexceeds 20%, strain and stress are locally concentrated, and thusductility is poor.

In addition, in an aspect of the present invention, an effect ofreducing the phase hardness difference between the soft phase and thehard phase can be obtained by containing a bainite phase in addition toa martensite phase as the hard phase.

The martensite phase and the bainite phase may be contained at 60% orless by area fraction, and the bainite phase may be contained at 10% ormore by area fraction.

The balance other than the soft phase and the hard phase may include aresidual austenite phase, and the residual austenite phase may becontained at such an extent that it does not affect the securement ofthe desired properties in the present invention. For example, theresidual austenite phase may be contained at 5% or less (including 0%)by area fraction.

In an aspect of the present invention, the above-mentioned structure,that is, the structure in which the soft phase and the hard phase areuniformly formed, can be obtained by C, Si, Al, Mo and Cr in theabove-mentioned alloy composition satisfying the following Expression 1and controlling the manufacturing conditions to be described below.

{(2×(Si+Al))+Mo+Cr}/C≥15   Expression 1

(Here, each element refers to a weight content.)

In Expression 1, Si and Al are ferrite-stabilizing elements and promoteferrite transformation, and Mo and Cr are elements contributing to theimprovement of hardenability. On the other hand, C is an elementcontributing to the formation of martensite by promoting C concentrationin untransformed austenite.

Thus, it is possible to obtain an effect of increasing the solidsolution concentration of Si, Al, Mo and Cr in ferrite and enhancing thestrength of ferrite by solid solution strengthening by controlling theratio of the elements that affect stabilization of ferrite, improvementsin hardenability, and the formation of martensite.

In the high-strength steel sheet of the present invention having theabove-mentioned structure, the concentration ratio of Si, Al, Mo, Cr andC in the ferrite satisfies the following Expression 2, so that the phasehardness difference, that is, the hardness ratio of a martensite phase,a bainite phase and a ferrite phase may satisfy the following Expression3.

{(2×(Si_(F)+Al_(F)))+Mo_(F)+Cr_(F)}/C_(F)≥500   Expression 2

(Here, each element refers to a weight content)

(H_(B)−H_(M))/(2×H_(F))≥3   Expression 3

(Here, B is bainite, M is martensite, and F is ferrite)

When the value of Expression 1 is less than 20, the effect of solidsolution strengthening by Si, Al, Mo and Cr cannot be sufficientlyobtained. Therefore, a concentration ratio (Expression 2) of Si, Al, Mo,Cr and C in the ferrite of 500 or more cannot be ensured. That is, thephase hardness difference is not effectively reduced, so that thehardness ratio H_(F) of ferrite which is a soft phase, and the hardnessratio of bainite H_(B) and martensite H_(M) which are hard phases,cannot be ensured to be 3 or less.

The contents of Si, Al, Mo, Cr and C in the steel, the contents of Si,Al, Mo, Cr and C in the ferrite and the hardness value of each phase maybe measured at the 1/4 t (Here, t is the thickness of the steel sheet)point in a thickness direction, but the present invention is not limitedthereto.

The high-strength steel sheet of the present invention has the structureas described above, so that the phase hardness difference may beminimized, and local stress concentration may be alleviated by finelydistributing each phase, and thereby ductility may be greatly improved.

Specifically, the high-strength steel sheet of the present invention mayhave a tensile strength of 980 MPa or more, a three-point bending anglemay be 80 degrees or more, and a product (YS×El) of yield strength andelongation may be 10000 or more.

Further, the high strength steel sheet of the present invention mayinclude a galvanized layer on at least one side.

Hereinafter, a method for manufacturing high-tensile steel havingexcellent processability according to another aspect of the presentinvention will be described in detail.

Briefly, the present invention can produce a target high-strength steelsheet through the processes of [reheating of steel slab-hotrolling-coiling-cold rolling-continuous annealing-cooling-hotgalvanizing-cooling], and the conditions of each step will be describedin detail as follows.

[Reheating of Steel Slab]

First, the steel slab having the above-mentioned component system isreheated. This step is performed in order to smoothly perform thesubsequent hot rolling step and sufficiently obtain the target physicalproperties of the steel sheet. In the present invention, the processcondition of the reheating process is not particularly limited, and maybe common conditions. As an example, a reheating process may beperformed in a temperature range of 1050 to 1300° C.

[Hot Rolling]

The steel slab heated as above may be finish hot rolled at a temperatureof the Ar3 transformation point or higher, and the temperature at theoutlet side preferably satisfies Ar3 to Ar3+50° C.

When the temperature at the outlet side during the finish hot rolling isless than Ar3, there may be a concern that performing ferrite andaustenite two-phase rolling leads to non-uniformity of the material. Onthe other hand, when the temperature exceeds Ar3+50° C., non-uniformityof the material may be caused due to the formation of abnormal coarsenedgrains due to high-temperature rolling, thereby causing a coil twistingphenomenon during subsequent cooling.

Meanwhile, the temperature at the inlet side may be in the range of 800to 1000° C. during the finish hot rolling.

[Coiling]

It is preferable to coiling the hot-rolled steel sheet produced asabove.

The coiling is preferably performed in the temperature range of 400 to700° C. When the coiling temperature is less than 400° C., an excessiveincrease in the strength of the hot-rolled steel sheet is caused due tothe formation of an excessive amount of martensite or bainite, and thusthere may arise a problem such as a shape defect due to a load in thesubsequent cold rolling. On the other hand, when the coiling temperatureexceeds 700° C., the surface concentration and internal oxidation ofelements which decrease the wettability of hot-dip galvanized steel suchas Si, Mn and B in the steel may be increased excessively.

[Cold Rolling]

The coiled hot-rolled steel sheet may be cold-rolled and produced as acold-rolled steel sheet.

In an aspect of the present invention, it is preferable that the coldrolling is performed at a total reduction ratio of 30% or more, and thereduction ratios of the first to fourth stands of the cold rollingstands are each set to be 15% or more, preferably. For example, thenumber of the cold rolling stands may be six.

This cold rolling increases the stored energy in the steel to act as adriving force for promoting the recrystallization of ferrite in thesubsequent annealing process, and may ultimately achieve the effect ofreducing the fraction of the non-recrystallized ferrite. When an excessamount of the non-recrystallized ferrite phase is present in the steel,strain and stress are locally concentrated, thereby reducing ductility.On the other hand, the recrystallized ferrite phase alleviates strainand stress concentration to contribute to ductility improvement.

When the total reduction ratio during the cold rolling is less than 30%,or the reduction ratio of each of the first to fourth stands is lessthan 15%, it is difficult to secure a target thickness and to correctthe shape of the steel sheet. Further, the fraction of thenon-recrystallized ferrite is formed to exceed 20% of the total fractionof the ferrite phase, resulting in poor ductility.

[Continuous Annealing]

It is preferable to continuously anneal the cold-rolled steel sheetproduced as above. The continuous annealing treatment may be performed,for example, in a continuous galvannealing line.

The continuous annealing step is a step for performingrecrystallization, forming ferrite and austenite phases and decomposingcarbon.

The continuous annealing treatment is preferably performed in thetemperature range of Ac1+30° C. to Ac3−20° C., and more preferably inthe temperature range of 780 to 820° C.

When the temperature is less than Ac1+30° C. during the continuousannealing, sufficient recrystallization cannot be achieved, and it isdifficult to form sufficient austenite so that the target level of themartensite phase and bainite phase fraction cannot be secured afterannealing. On the other hand, when the temperature exceeds Ac3−20° C.,productivity is lowered, the austenite phase is excessively formed, thefractions of martensite and bainite phases are greatly increased aftercooling, and thus it becomes difficult to secure desired ductility.Further, the surface concentration due to the elements such as Si, Mn, Band the like, which deteriorate the wettability of the hot-dipgalvanizing becomes severe, and thus the quality of the plated surfacemay be lowered.

[Stepwise Cooling]

It is preferable that the cold-rolled steel sheet subjected to thecontinuous annealing treatment as above is cooled step-by-step.

Specifically, it is preferable that the cooling is primarily cooled upto 630 to 670° C. at an average cooling rate of 10° C./s or less(excluding 0° C./s), and then secondarily cooled up to 400 to 550° C. atan average cooling rate of 5° C./s or more.

When the stop temperature of the primary cooling is less than 630° C.,the diffusion activity of carbon is low due to an excessively lowtemperature, and thus the concentration of carbon in ferrite becomeshigh and the yield ratio increases. Accordingly, the tendency ofoccurrence of cracks during processing increases. On the other hand,when the stop temperature of the primary cooling is higher than 670° C.,it is advantageous in terms of diffusion of carbon, but it isdisadvantageous that an excessively high cooling rate is required in thesubsequent cooling (secondary cooling). In addition, when the averagecooling rate during the primary cooling exceeds 10° C./s, the diffusionof carbon cannot sufficiently occur. Meanwhile, the lower limit of theaverage cooling rate is not particularly limited, but may be 1° C./s ormore in consideration of productivity.

It is preferable to carry out the secondary cooling after completion ofthe primary cooling under the above conditions. When the stoptemperature of the secondary cooling exceeds 550° C., the bainite phasecannot be sufficiently secured. Meanwhile, when the stop temperature ofthe secondary cooling is less than 400° C., the fraction of themartensite phase becomes excessive, and thus it is difficult to securethe desired ductility. Further, when the average cooling rate during thesecondary cooling is less than 5° C./s, the bainite phase may not beformed at the target level. On the other hand, the upper limit of theaverage cooling rate is not particularly limited, and may be suitablyselected by those skilled in the art in consideration of thespecifications of the cooling equipment. For example, it may beperformed at 100° C./s or lower.

In third cooling, hydrogen cooling equipment using hydrogen gas (H₂ gas)may be used. In this way, when cooling is performed using the hydrogencooling equipment, it is possible to obtain an effect of suppressingsurface oxidation that may occur in the tertiary cooling.

Meanwhile, in the stepwise cooling as above, the cooling rate at thetime of the secondary cooling may be faster than the cooling rate at thetime of the primary cooling.

[Maintaining]

It is preferable that the temperature is maintained for 50 to 500seconds in the cooled temperature range after completing the stepwisecooling as described above.

When the maintaining time is less than 50 seconds, the bainite phase maynot be sufficiently formed. On the other hand, when the maintaining timeexceeds 500 seconds, an excess amount of the bainite phase is formed andthus it may be difficult to secure the target microstructure.

[Hot-Dip Galvanizing]

It is preferable that the steel sheet is immersed in a hot-dipgalvanizing bath after the stepwise cooling and maintaining processes asabove to produce a hot-dip galvanized steel sheet.

Here, although the hot-dip galvanizing may be carried out under commonconditions, it may be carried out in the temperature range of 430 to490° C., for example. The composition of the hot-dip galvanizing bathduring the hot-dip galvanizing is not particularly limited, and may be apure galvanizing bath or a zinc-based alloy plating bath containing Si,Al, Mg, or the like.

[Final Cooling]

After the completion of the hot-dip galvanizing, it is preferable tocool the steel sheet at a cooling rate of 1° C./s or more to amartensitic transformation start temperature (Ms) or lower. In thisprocess, the martensite phase and the residual austenite phase may beformed in the steel sheet (the steel sheet corresponds to the basematerial of the lower part of the plated layer).

When the stop temperature of the cooling exceeds Ms, the martensitephase cannot be sufficiently secured. When the average cooling rate isless than 1° C./s, the martensite phase is unevenly formed due to anexcessively slow cooling rate. More preferably, the cooling may beperformed at a cooling rate of 1 to 100° C./s.

Even when the steel sheet is cooled to room temperature during thecooling, there is no problem in securing the target structure, and theroom temperature may be expressed as about 10 to 35° C.

Meanwhile, if necessary, the hot-dip galvanized steel sheet may besubjected to alloying heat treatment before the final cooling to obtaina hot-dip galvannealed steel sheet. In the present invention, theprocess conditions of the alloying heat treatment are not particularlylimited, and may be common conditions. As an example, the alloying heattreatment may be performed at a temperature of 480 to 600° C.

Next, if necessary, the hot-dip galvanized steel sheet or the hot-dipgalvannealed steel sheet finally cooled is subject to temper rolling toform large amounts of dislocations in ferrite disposed aroundmartensite, thereby further improving bake hardenability.

Here, a reduction ratio is preferably less than 1.0% (excluding 0%).When the reduction ratio is 1.0% or more, it is advantageous in terms offormation of dislocation, but it may cause side effects such asoccurrence of plate breakage due to an equipment capability limit.

The high-strength steel sheet of the present invention produced as aboveincludes a mixture of a hard phase and a soft phase as a microstructure,and specifically may include ferrite having an area fraction of lessthan 40%, and the remainder as bainite, martensite and residualaustenite.

Here, the concentration of the solid solution elements in the ferritemay be increased to improve strength and implement grain refinement,thereby minimizing the difference in hardness between the soft phase andthe hard phase. Further, there is an effect of providing a high-strengthsteel sheet having excellent bending properties and formability byimproving ductility in addition to high yield strength.

Hereinafter, the present disclosure will be described more specificallyaccording to examples. However, the following examples should beconsidered in a descriptive sense only and not for purposes oflimitation. The scope of the present invention is defined by theappended claims, and modifications and variations may reasonably madetherefrom.

EXAMPLES

After a steel slab having the alloy composition shown in the followingTable 1 was prepared, the steel slab was heated to a temperature rangeof 1050 to 1250° C., and then subjected to hot rolling, cooling andcoiling under the conditions shown in the following Table 2 to prepare ahot rolled steel sheet.

Thereafter, each hot-rolled steel sheet was cold-rolled to produce acold-rolled steel sheet, which was subjected to continuous annealingtreatment under the conditions shown in the following Table 2, followedby stepwise cooling (primary and secondary) and maintained at thesecondary cooling stop temperature for 50 to 500 seconds. Here, thesecondary cooling was performed using hydrogen cooling equipment.

Thereafter, the steel sheet was galvanized at 430 to 490° C. in ahot-dip galvanizing bath (0.1 to 0.3% Al-remainder Zn), and thensubjected to final cooling at a cooling rate of 1° C./s or more up to Msor less and subjected to temper rolling at 0.2% to produce a hot-dipgalvanized steel sheet.

The microstructure of each steel sheet manufactured above was observed,and mechanical properties and plating characteristics were evaluated.The results are shown in the following Table 3.

Here, the tensile test for each test piece was performed in the Ldirection using the ASTM standard. The bending angle (180degrees-bending angle) was evaluated by applying the German Associationof the Automotive Industry (VDA) 238-100 standard for the three-pointbending test. The larger the bending angle, the more excellent thebending property.

Further, the microstructure fraction was assessed by analyzing a matrixat a point of 1/4 t of the thickness of the steel sheet. Specifically,the fractions of ferrite, bainite, martensite and austenite weremeasured by FE-SEM and an image analyzer after Nital corrosion.

The concentrations of C, Si, Al, Mo, and Cr in the ferrite at a point of¼ t of each steel sheet were measured using transmission electronmicroscopy (TEM), energy dispersive spectroscopy (EDS), and ELLSanalysis equipment.

Further, the hardness of each phase was measured ten times using a MicroVickers hardness tester, and then an average value was obtained.

TABLE 1 Alloy composition (wt %) Composition Classification C Si Mn P SAl Mo Cr Ti Nb N B Sb ratio Inventive 0.07 0.4 2.35 0.02 0.003 0.04 0.120.85 0.02 0.025 0.005 0.0025 0.03 26.4 steel 1 Inventive 0.09 0.3 2.40.01 0.005 0.40 0.10 0.65 0.022 0.034 0.003 0.0003 0.03 23.9 steel 2Inventive 0.08 0.5 2.4 0.009 0.004 0.06 0.09 0.71 0.04 0.03 0.004 0.00250.01 24.0 steel 3 Inventive 0.10 0.6 2.3 0.01 0.003 0.07 0.12 0.68 0.030.02 0.0056 0.0025 0.02 21.4 steel 4 Inventive 0.06 0.4 2.2 0.01 0.0030.05 0.15 0.65 0.02 0.03 0.005 0.0011 0.03 28.3 steel 5 Inventive 0.070.4 2.2 0.01 0.003 0.04 0.07 0.60 0.025 0.02 0.0047 0.0005 0.02 22.1steel 6 Inventive 0.13 0.6 2.2 0.01 0.003 0.06 0.13 0.71 0.003 0.020.005 0.0020 0.02 16.6 steel 7 Comparative 0.10 0.1 2.5 0.02 0.002 0.0350.07 0.95 0.02 0.05 0.005 0.0025 0.02 12.9 steel 1 Comparative 0.06 0.42.9 0.02 0.002 0.03 0.01 0.02 0.03 0.02 0.005 0.0010 0.01 14.8 steel 2Comparative 0.07 0.2 2.5 0.02 0.002 0.30 0.04 0.04 0.015 0.02 0.0050.0014 0.01 15.4 steel 3 Comparative 0.09 0.2 2.7 0.01 0.003 0.10 0.010.60 0.03 0.02 0.0056 0.0025 0.02 13.4 steel 4 Comparative 0.15 0.3 2.50.01 0.003 0.04 0.07 0.40 0.025 0.02 0.0047 0.0005 0.02 7.7 steel 5

(In Table 1, the composition ratios are the values of Expression1[{(2×(Si+Al))+Mo+Cr}/C] of each steel.)

TABLE 2 Cold rolling (%) Temperature Coiling Total Reduction ratioAnnealing Primary cooling Secondary cooling at outlet side temperaturereduction of First to temperature Rate Temperature Rate TemperatureClassification (° C.) (° C.) ratio forth stands (° C.) (° C./s) (° C./s)(° C./s) (° C./s) Inventive 920 650 70 15 or more 810 3.0 630 7.1 520steel 1 Inventive 911 600 68 15 or more 790 2.9 652 8.9 530 steel 2Inventive 907 580 61 15 or more 820 4.0 649 10.7 500 steel 3 Inventive905 629 80 15 or more 790 3.8 664 13.2 480 steel 4 Inventive 915 647 7115 or more 800 5.2 645 15.1 420 steel 5 Inventive 900 680 64 15 or more810 7.1 670 12.5 480 steel 6 Inventive 905 630 78 15 or more 800 4.0 65012.0 450 steel 7 Comparative 800 400 63 15 or more 780 10.4 650 5.1 580steel 1 Comparative 830 650 65 15 or more 790 4.4 590 5.0 520 steel 2Comparative 900 716 28 10 or less 810 5.7 630 7.2 520 steel 3Comparative 916 615 71 15 or more 820 2.0 630 10.2 450 steel 4Comparative 770 350 25 10 or less 850 8.1 530 15.1 300 steel 5

TABLE 3 Mechanical properties Three-point Microstructure (%) YS TS EI YS× EI bending angle Hardness Concentration Classification F B + M Fn/Ft(MPa) (MPa) (%) YR (MPa %) (degree) ratio ratio Inventive 34 66 0 7191040 15 0.69 10785 91 2.6 617 steel 1 Inventive 31 69 0 751 1038 15 0.7211265 95 2.8 717 steel 2 Inventive 35 65 0 748 1039 16 0.72 11968 87 2.5640 steel 3 Inventive 37 63 0 701 1016 17 0.69 11917 95 2.4 713 steel 4Inventive 29 71 0 718 1017 15 0.71 10770 89 2.1 567 steel 5 Inventive 3367 0 736 1055 15 0.70 11040 85 2.8 517 steel 6 Inventive 36 64 0 7351027 16 0.72 11760 89 2.4 670 steel 7 Comparative 43 57 0 640 1070 150.62 9600 67 3.1 430 steel 1 Comparative 46 54 0 666 999 13 0.67 8658 733.4 297 steel 2 Comparative 35 65 21 750 1005 12 0.71 9000 78 3.3 360steel 3 Comparative 45 55 0 699 1014 12 0.69 8388 69 4.1 403 steel 4Comparative 24 76 34 838 1105 11 0.76 9218 63 3.7 383 steel 5

(In Table 3, F is a ferrite phase, B is a bainite phase, and M is amartensite phase. In addition, YS is yield strength, TS is tensilestrength, El is elongation, and YR is yield ratio. The concentrationratio is calculated by Expression2[{(2×(Si_(F)+Al_(F)))+Mo_(F)+Cr_(F)}/C_(F)], and the hardness ratio iscalculated by Expression 3[(H_(B)+H_(M))/(2×H_(F))]).

As shown in Tables 1 to 3, since the intended microstructure was formed,inventive steels 1 to 7, in which the steel alloy composition, thecomposition ratio (Expression 1), and the manufacturing conditionssatisfy all the requirements of the present invention, have athree-point bending angle of 80 degrees or more, and a product of yieldstrength and elongation (YS×El) of 10000 or more, and thuscrashworthiness and formability can be secured.

On the other hand, comparative steels 1 to 5, in which one or more ofthe conditions of the steel alloy composition, the composition ratio andthe manufacturing conditions deviate from those proposed in the presentinvention, fail to obtain the intended microstructure in the presentinvention, bending properties are deteriorated, and a product of yieldstrength and elongation (YS×El) of less than 10000 was secured, so thatcrashworthiness and formability could not be ensured.

FIG. 1 shows a change in the interphase hardness ratio[(H_(B)+H_(M))/(2×H_(F))] according to a concentration ratio(corresponding to Expression 2) of Si, Al, Mo, Cr and C in a ferritephase.

As shown in FIG. 1, it can be seen that when a concentration ratiobetween Si, Al, Mo, Cr, and C in the ferrite phase of 500 or more issecured, the difference in hardness between phases can be effectivelyreduced.

FIG. 2 shows a change in the product of yield strength and elongation(YS×El) according to the interphase hardness ratio[(H_(B)+H_(M))/(2×H_(F))].

As shown in FIG. 2, it can be seen that when the interphase hardnessratio is 3 or less, a product of yield strength and elongation (YS×El)of 10000 or more can be secured.

FIG. 3 shows a change in a three-point bending angle according to theinterphase hardness ratio [(H_(B)+H_(M))/(2×H_(F))].

As shown in FIG. 3, it can be seen that when the interphase hardnessratio is 3 or less, a three-point bending angle of 80 degree or more canbe secured.

FIG. 4 shows a change in a three-point bending angle according to theproduct of yield strength and elongation (YS×El).

As shown in FIG. 4, it can be seen that when the value of therelationship between the yield strength and elongation (YS×El) is 10000or more, a three-point bending angle of 80 degrees or more can besecured.

What is claimed are:
 1. A high-strength steel sheet comprising, in %weight, carbon (C): 0.04 to 0.15%, silicon (Si): 0.01 to 1.0%, manganese(Mn): 1.8 to 2.5%, molybdenum (Mo): 0.15% or less (excluding 0%),chromium (Cr): 1.0% or less (excluding 0%), phosphorus (P): 0.1% orless, sulfur (S): 0.01% or less, aluminum (Al): 0.01 to 0.5%, nitrogen(N): 0.01% or less, boron (B): 0.01% or less (excluding 0%), antimony(Sb):
 0. 05% or less (excluding 0%), one or more of titanium (Ti): 0.003to 0.06% and niobium (Nb): 0.003 to 0.06%, a balance of Fe and otherunavoidable impurities, and contents of the C, the Si, the Al, the Moand the Cr satisfying the following Expression 1:{(2×(Si+Al))+Mo+Cr}/C≥15   Expression 1 wherein the high-strength steelsheet comprises: a microstructure including a ferrite phase, a bainitephase, a martensite phase, and a residual austenite phase, the ferritephase being less than 40% of area fraction in the microstructure, andwherein an area ratio (Fn/Ft) of non-recrystallized ferrite (Fn) in theferrite phase (Ft) is 20% or less (including 0%).
 2. The high-strengthsteel sheet of claim 1, wherein contents of the Si, the Al, the Mo, theCr and the C in the ferrite phase satisfy the following Expression 2:{(2×(Si_(F)−Al_(F)))+Mo_(F)+Cr_(F)}/C_(F)≥500   Expression 2 (whereSi_(F) refers to a content of Si, Al_(F) refers to a content of Al,Mo_(F) refers to a content of Mo, and Cr_(F) refers to a content of Cr).3. The high-strength steel sheet of claim 1, wherein a hardness ratio ofthe martensite phase, the bainite phase and the ferrite phase satisfiesthe following Expression 3:(H_(B)−H_(M))/(2×H_(F))≥3   Expression 3 (where H_(B) refers to ahardness of bainite phase, H_(M) refers to a hardness of martensitephase, and H_(F) refers to a hardness of ferrite phase).
 4. Thehigh-strength steel sheet of claim 1, wherein the high-strength steelsheet further comprises: a hot dipped galvanized layer formed on atleast one surface thereof.
 5. The high-strength steel sheet of claim 1,wherein the high-strength steel sheet has a tensile strength of 980 MPaor more, a three-point bending angle is 80 degrees or more, and aproduct (YS×El) of yield strength and elongation is 10000 or more.
 6. Ahigh-strength steel sheet comprising, in % weight, carbon (C): 0.04 to0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.8 to 2.5%,molybdenum (Mo): 0.15% or less (excluding 0%), chromium (Cr): 1.0% orless (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.01% orless, aluminum (Al): 0.01 to 0.5%, nitrogen (N): 0.01% or less, boron(B): 0.01% or less (excluding 0%), antimony (Sb):
 0. 05% or less(excluding 0%), one or more of titanium (Ti): 0.003 to 0.06% and niobium(Nb): 0.003 to 0.06%, a balance of Fe and other unavoidable impurities,and contents of the C, the Si, the Al, the Mo and the Cr satisfying thefollowing Expression 1:{(2×(Si+Al))+Mo+Cr}/C≥15   Expression 1 wherein the high-strength steelsheet comprises: a microstructure including a ferrite phase, a bainitephase, a martensite phase, and a residual austenite phase, the ferritephase being less than 40% of area fraction in the microstructure, andwherein contents of the Si, the Al, the Mo, the Cr and the C in theferrite phase satisfy the following Expression 2:{(2×(Si_(F)−Al_(F)))+Mo_(F)+Cr_(F)}/C_(F)≥500   Expression 2 (whereSi_(F) refers to a content of Si, Al_(F) refers to a content of Al,Mo_(F) refers to a content of Mo, and Cr_(F) refers to a content of Cr).7. The high-strength steel sheet of claim 6, wherein a hardness ratio ofthe martensite phase, the bainite phase and the ferrite phase satisfiesthe following Expression 3:(H_(B)−H_(M))/(2×H_(F))≥3   Expression 3 (where H_(B) refers to ahardness of bainite phase, H_(M) refers to a hardness of martensitephase, and H_(F) refers to a hardness of ferrite phase).
 8. Thehigh-strength steel sheet of claim 6, wherein the high-strength steelsheet further comprises: a hot dipped galvanized layer formed on atleast one surface thereof.
 9. The high-strength steel sheet of claim 6,wherein the high-strength steel sheet has a tensile strength of 980 MPaor more, a three-point bending angle is 80 degrees or more, and aproduct (YS×El) of yield strength and elongation is 10000 or more.
 10. Ahigh-strength steel sheet comprising, in % weight, carbon (C): 0.04 to0.15%, silicon (Si): 0.01 to 1.0%, manganese (Mn): 1.8 to 2.5%,molybdenum (Mo): 0.15% or less (excluding 0%), chromium (Cr): 1.0% orless (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.01% orless, aluminum (Al): 0.01 to 0.5%, nitrogen (N): 0.01% or less, boron(B): 0.01% or less (excluding 0%), antimony (Sb):
 0. 05% or less(excluding 0%), one or more of titanium (Ti): 0.003 to 0.06% and niobium(Nb): 0.003 to 0.06%, a balance of Fe and other unavoidable impurities,and contents of the C, the Si, the Al, the Mo and the Cr satisfying thefollowing Expression 1:{(2×(Si+Al))+Mo+Cr}/C≥15   Expression 1 wherein the high-strength steelsheet comprises: a microstructure including a ferrite phase, a bainitephase, a martensite phase, and a residual austenite phase, the ferritephase being less than 40% of area fraction in the microstructure, andwherein a hardness ratio of the martensite phase, the bainite phase andthe ferrite phase satisfies the following Expression 3:(H_(B)−H_(M))/(2×H_(F))≥3   Expression 3 (where H_(B) refers to ahardness of bainite phase, H_(M) refers to a hardness of martensitephase, and H_(F) refers to a hardness of ferrite phase).
 11. Thehigh-strength steel sheet of claim 10, wherein an area ratio (Fn/Ft) ofnon-recrystallized ferrite (Fn) in the ferrite phase (Ft) is 20% or less(including 0%), and wherein contents of the Si, the Al, the Mo, the Crand the C in the ferrite phase satisfy the following Expression 2:{(2×(Si_(F)−Al_(F)))+Mo_(F)+Cr_(F)}/C_(F)≥500   Expression 2 (whereSi_(F) refers to a content of Si, Al_(F) refers to a content of Al,Mo_(F) refers to a content of Mo, and Cr_(F) refers to a content of Cr).12. The high-strength steel sheet of claim 10, wherein the high-strengthsteel sheet further comprises: a hot dipped galvanized layer formed onat least one surface thereof.
 13. The high-strength steel sheet of claim10, wherein the high-strength steel sheet has a tensile strength of 980MPa or more, a three-point bending angle is 80 degrees or more, and aproduct (YS×El) of yield strength and elongation is 10000 or more.